In recent years, in the automobile industry, high-strength steel sheet has been used to achieve a combination of functions for protecting the occupants in the case of a collision and a reduction in weight that improves fuel consumption. In terms of ensuring favorable safety during a collision, heightened appreciation of safety factors and more stringent regulations mean that there is now a need to use high-strength steel sheet for components of complex shape, which until now have been manufactured using low-EM strength steel sheet. For this reason, superior hole expansion properties are now being demanded for high-strength steel.
Many components within an automobile are joined using welding techniques such as spot welding, arc welding or laser welding, and therefore in order to enhance the collision safety for the vehicle, it is necessary that these joins do not fracture upon collision. In other words, if a fracture occurs at a joint upon collision, then even if the strength of the steel is adequate, the joint structure is unable to satisfactorily absorb the energy of the collision, making it impossible to achieve the required collision energy absorption performance.
Accordingly, automobile components must also exhibit excellent joint strength for joints manufactured by spot welding, arc welding, laser welding, or the like. However, a problem arises in that as the amounts of C, Si, Mn, and the like are increased to achieve greater strength for the steel sheet, an accompanying deterioration in the strength of the welded portions tends to occur, meaning it is desirable that strengthening of the steel is achieved without excessive increases in the amounts of the alloy elements incorporated within the steel.
Examples of indicators for evaluating the strength of a spot welded joint include a tensile shear strength (TSS) test prescribed in JIS Z 3136 in which a shear stress is applied to the weld, and a cross tension strength (CTS) test prescribed in JIS Z 3137 in which stress is applied in the direction of joint separation. Of these two tests, it is known that the TSS value increases with increasing steel sheet strength, whereas the CTS value does not increase even with an increase in the steel sheet strength. As a result, the ductility ratio, which is represented by the ratio between TSS and CTS, decreases with increased addition of alloy components to the steel, namely with increased steel strength. It is well known that high-strength steel sheet having a high C content has problems in terms of spot weldability (see Non-Patent Document 1).
On the other hand, formability of a material tends to deteriorate as the strength of the material is increased, and if a high-strength steel sheet is to be used for forming a member having a complex shape, then a steel sheet that satisfies both of favorable formability and high strength must be manufactured. Although the simple term “formability” is used, when applied to a member having a complex shape such as an automobile component, the component actually requires a combination of a variety of different formability properties including ductility, stretch formability, bendability, hole expandability, and stretch flange formability.
It is known that the ductility and the stretch formability correlate with the work hardening coefficient (the n value), and steel sheets having high n values are known to exhibit excellent formability. Examples of steel sheets that exhibit excellent ductility and stretch formability include DP (Dual Phase) steel sheets in which the microstructure of the steel sheet is composed of ferrite and martensite, and TRIP (Transformation Induced Plasticity) steel sheets in which the microstructure of the steel sheet includes residual austenite.
On the other hand, known examples of steel sheets that exhibit excellent hole expandability include steel sheets having a precipitation-strengthened ferrite single phase microstructure, and steel sheets having a bainite single phase microstructure (see Patent Documents 1 to 3, and Non-Patent Document 2).
Further, it is known that the bendability correlates with the structural uniformity, and it has been demonstrated that the bendability can be improved by improving the uniformity of the steel microstructure (see Non-Patent Document 3).
Accordingly, steel sheets in which the steel microstructure is formed as a precipitation-strengthened ferrite single phase microstructure (Non-Patent Document 2) and DP steel sheets which, although having dual phase microstructures composed of ferrite and martensite, exhibit enhanced uniformity as a result of miniaturization of the steel microstructures (see Patent Document 4) are already known.
DP steel sheets contain highly ductile ferrite as the main phase, and by dispersing martensite which is the hard microstructure within the microstructure of the steel sheet, excellent ductility can be achieved. Furthermore, the softer ferrite is easily molded, and because a large amount of dislocation is introduced at the same time as the molding, and is subsequently hardened, the n value is high. However, if the steel microstructure is composed of soft ferrite and hard martensite, then because the molding capabilities of the two microstructures differ, when molding is conducted as part of large scale operations such as hole expansion processing, minute microvoids tend to form at the interfaces between the two different microstructures, resulting in a marked deterioration in the hole expandability. The volume fraction of martensite incorporated within the DP steel sheet having a maximum tensile strength of 590 MPa or higher is comparatively large, and because the steel also contains a multitude of ferrite-martensite interfaces, the microvoids formed at these interfaces can readily interconnect, which can lead to cracking and fracture. For these reasons, the hole expandability properties of the DP steel sheets is poor (see Non-Patent Document 4).
It is known that a microstructure containing tempered martensite can be used to improve the hole expandability in these DP steel sheets composed of ferrite and martensite (see Patent Document 5). However, it is necessary to conduct an additional tempering treatment in order to improve the hole expandability; therefore, productivity problems arise. Moreover, a decrease in the strength of the steel sheet due to the tempered martensite is also unavoidable. As a result, the amount of C added to the steel must be increased to maintain the strength of the steel, but this causes a deterioration in the weldability. In other words, with regard to the DP steel sheets formed from ferrite and martensite, achieving both strength in the order of 880 MPa, as well as favorable hole expandability and weldability has proven impossible.
In addition, when tempered martensite is converted to a hard microstructure, the volume fraction of ferrite must be reduced in order to maintain the strength; however, this results in a deterioration in the ductility.
Furthermore, in a development related to the DP steel sheet, a high-tensile hot-dip galvanized steel sheet has been proposed that is composed of ferrite and a hard second phase, and this steel exhibits excellent balance between strength and ductility, as well as superior balance between bendability, spot weldability, and plating adhesion (see Patent Document 6). As the hard second phase, martensite, bainite, and residual austenite are exemplified. However, with regard to this high-tensile hot-dip galvanized steel sheet, annealing must be conducted at a high temperature within a range from A3 to 950° C.; therefore, there is a problem that the productivity is poor. In particular, if achieving favorable spot weldability is also taken into consideration, then the amount of C, which functions as an austenite stabilizing element (namely, an element that lowers the Ac3 point) added to the steel must be suppressed, which frequently results in high annealing temperatures and reduced productivity. Moreover, annealing at extremely high temperatures exceeding 900° C. is undesirable, because it can cause severe damage to the production equipments such as the furnace casing and the hearth roll, and it tends to promote the formation of surface defects on the surface of the steel sheet.
Further, with regard to the high-tensile hot-dip galvanized steel sheet proposed in Patent Document 6, the hole expandability is 55% at 918 MPa, 35% at 1035 MPa, 35% at 1123 MPa, and approximately 26% at 1253 MPa. In comparison, the hole expandability results for the present invention are 90% at 980 MPa, 50% at 1080 MPa, and 40% at 1180 MPa, indicating that with regard to the high-tensile hot-dip galvanized steel sheet of Patent Document 6, it impossible to achieve a satisfactory combination of strength and hole expandability.
The hole expandability ends to be similarly low in TRIP steel sheets in which the steel microstructure is composed of ferrite and residual austenite. This is because mold working of automobile components, including hole expanding and stretch flange forming, is conducted after punching out or mechanical cutting of the sheet.
The residual austenite contained within the TRIP steel sheets transforms into martensite when subjected to processing. For example, drawing or stretching of the steel causes the residual austenite to transform into martensite; thereby, increasing the strength of the processed portions, and by restricting the concentration of this transformation, a high degree of formability can be maintained.
However, when the steel is punched out or cut, the portions close to the edges are subjected to processing, and therefore the residual austenite incorporated within the steel microstructure in these portions transforms into martensite. As a result, a microstructure similar to that of a DP steel sheet is obtained, and the hole expandability and stretch flange formability tend to deteriorate. Alternatively, because the punching out process itself is a process that accompanies large deformation, it has been reported that after punching out of the steel, microvoids tend to exist at the interfaces between the ferrite and hard microstructures (in this case, the martensite formed by transformation of the residual austenite), resulting in a deterioration in the hole expandability. Moreover, steel sheets in which cementite or pearlite microstructures exist at the grain boundaries also exhibit poor hole expandability. This is because the interfaces between the ferrite and cementite act as origins for microscopic void formation.
Furthermore, in order to ensure that the residual austenite is maintained, a large amount of C must be concentrated within the austenite; however, compared with a DP steel having the same C content (a multi-phase steel sheet composed of ferrite and martensite), the volume fraction of hard microstructures tends to decrease, making it difficult to maintain strength. In other words, in the case in which a high strength of at least 880 MPa is ensured, the amount of added C required for strengthening increases considerably; thereby, causing a deterioration in the spot weldability. Accordingly, the upper limit for the volume fraction of residual austenite is 3%.
As a result, as disclosed in Patent Documents 1 to 3, research into steel sheets having excellent hole expandability has led to the development of high-strength hot-rolled steel sheets having single phase microstructure of either bainite or precipitation-strengthened ferrite as the main phase, in which a large amount of an alloy-carbide-forming element such as Ti is added to convert the C incorporated within the steel into an alloy carbide; thereby, suppressing the formation of a cementite phase at the grain boundaries, and yielding superior hole expandability.
In the case of a steel sheet having a bainite single phase microstructure, in order to convert the microstructure of the steel sheet to a bainite single phase microstructure, the production of the cold-rolled steel sheet must include first heating to a high temperature to form an austenite single phase; therefore, the productivity is poor. Furthermore, bainite microstructures include a large amount of dislocation; therefore, they exhibit poor workability and are difficult to use for components that require favorable ductility and stretch formability. Furthermore, if consideration is given to ensuring a high strength of at least 880 MPa, then an amount of C exceeding 0.1% by mass must be added, which means the steel suffers the aforementioned problem of being unable to achieve a combination of high strength and favorable spot weldability.
In steel sheets having a precipitation-strengthened ferrite single phase microstructure, precipitation strengthening provided by carbides of Ti, Nb, Mo, V, or the like is used to increase the strength of the steel sheet while suppressing the formation of cementite and the like; thereby, a steel sheet having a combination of a high strength of 880 MPa or higher and superior hole expandability can be obtained. However, in the case of cold-rolled steel sheets that undergo cold rolling and annealing steps, it is difficult to utilize the above precipitation strengthening effect.
In other words, the precipitation strengthening is accomplished by coherent precipitation of an alloy carbide of Nb or Ti or the like within the ferrite. In a cold-rolled steel sheet that has been subjected to cold rolling and annealing, because the ferrite is processed and is recrystallized during annealing, the orientation relationship with the coherent precipitated Nb or Ti precipitate during the hot rolling stage is lost; therefore, the strengthening function of the precipitate is largely lost, and making it difficult to use this technique for strengthening cold-rolled steel.
Further, it is known that when cold rolling is conducted, the Nb or Ti significantly delay the recrystallization, meaning that in order to ensure excellent ductility, a high-temperature annealing step is required, which results, in poor productivity. Furthermore, even if ductility similar to that of hot-rolled steel sheet were to be obtained, precipitation-strengthened steel still exhibits inferior ductility and stretch formability; therefore, it is unsuitable for regions that require superior stretch formability.
Here, in the present invention, a steel sheet of which the product of the maximum tensile strength and the total elongation is 16,000 (MPa×%) or more is deemed to be high-strength steel having favorable ductility. In other words, the targeted ductility values are 18.2% at 880 MPa, 16.3% or greater at 980 MPa, 14.8% or greater at 1080 MPa, and 13.6% or greater at 1180 MPa.
Steel sheets that address these problems and are provided to satisfy a combination of superior ductility and hole expandability are disclosed in Patent Documents 7 and 8. These steel sheets are manufactured by initially forming a multi-phase microstructure composed of ferrite and martensite, and subsequently tempering and softening the martensite; thereby, an attempt is made to yield an improved balance between the strength and ductility, as well as a simultaneous improvement in the hole expandability, by structurally strengthening the steel.
However, even if improvements in the hole expandability and stretch flange formability are achieved by softening of hard microstructures due to tempering of the martensite, the problem of inferior spot weldability remains if applied to high-strength steel sheets of 880 MPa or higher.
For example, by tempering martensite, hard microstructures can be softened and the hole expandability can be improved. However, because a reduction in the strength also occurs simultaneously, the volume fraction of martensite must be increased so as to offset this reduction in strength; therefore, a large amount of C must be added. As a result, spot weldability and the like tend to deteriorate. Furthermore, in the case of using equipments such as hot-dip galvanizing equipment in which both of quenching and tempering cannot be conducted, a microstructure containing ferrite and martensite microstructure must first be formed, and a separate heat treatment must then be conducted; therefore, the productivity is poor.
On the other hand, it is well known that the strength of a welded joint is dependent on the amount of added elements, and particularly added C, contained within the steel sheet. It is known that by strengthening a steel sheet while restricting the amount of C added, a combination of favorable strength and favorable weldability (namely, maintenance of the joint strength of a welded portion) can be obtained. Because a welded portion is melted and then cooled at a rapid cooling rate, the microstructure of the hard portion becomes to mainly include martensite. Accordingly, the welded portion is extremely hard and exhibits poor deformability (molding capabilities). Moreover, even if the microstructure of the steel sheet has been controlled, because the steel is melted upon welding, control of the microstructure within the welded portion is extremely difficult. As a result, improvements in the properties of the welded portion have conventionally been made by controlling the components within the steel sheet (for example, see Patent Document 4 and Patent Document 9).
The description above also applies to steel sheets having a multi-phase microstructure containing ferrite and bainite. In other words, a bainite microstructure is formed at a higher temperature than a martensite microstructure, and is therefore considerably softer than martensite. As a result, bainite microstructures are known to exhibit superior hole expandability. However, since they are soft microstructures, it is difficult to achieve a high strength of 880 MPa or higher. In those cases where the main phase is ferrite and the hard microstructures are formed as bainite microstructures, in order to ensure a high strength of at least 880 MPa, the amount of added C must be increased, the proportion of bainite microstructures must be increased, and the strength of the bainite microstructures must be improved. This causes a marked deterioration in the spot weldability of the steel.
Patent Document 9 discloses that by adding Mo to a steel sheet, favorable spot weldability properties can be achieved even for steel sheets having a C content exceeding 0.1% by mass. However, although adding Mo to the steel sheet suppresses the formation of voids or cracks within the spot welded portion, and improves the strength of the welded joint for welding conditions where these types of defects occur readily, there is no improvement in the strength of the welded joint under conditions where the above defects do not occur. Furthermore, if consideration is given to achieving a high strength of at least 880 MPa, then addition of a large amount of C is unavoidable, and the problem remains that it is difficult to obtain a steel sheet that exhibits both favorable spot weldability and superior formability. Furthermore, because the steel sheet includes residual austenite as the hard microstructure, during hole expansion or stretch flange formation, stress tends to be concentrated at the interfaces between the soft ferrite that represents the main phase and the residual austenite that functions as the hard microstructure, resulting in microvoid formation and interconnection; thereby, deterioration occurs in these properties.
Furthermore, Mo tends to promote the formation of band-like microstructures, causing a deterioration in the hole expandability. Accordingly, in the present invention, as described below, investigations were focused on conditions that realized satisfactory weldability without the addition of Mo.
A known steel sheet that combines a high maximum tensile strength of at least 780 MPa with favorable spot weldability is disclosed in Patent Document 4 listed below. In this steel sheet, by utilizing a combination of precipitation strengthening due to the addition of Nb or Ti, fine-grain strengthening, and dislocation strengthening that utilizes non-recrystallized ferrite, a steel sheet that combines a strength of at least 780 MPa with superior ductility and bendability can be obtained even when the carbon content of the steel sheet is 0.1% by mass or less. However, in order to enable application to components having more complex shapes, further improvements in the ductility and hole expandability are still required. As described above, achieving a combination of high strength of at least 880 MPa and superior levels of ductility, stretch formability, bendability, hole expandability, stretch flange formability, and spot weldability has proven extremely difficult.    Patent Document 1: Japanese Unexamined Patent Application, First Publication No. 2003-321733    Patent Document 2: Japanese Unexamined Patent Application, First Publication No. 2004-256906    Patent Document 3: Japanese Unexamined Patent Application, First Publication No. H11-279691    Patent Document 4: Japanese Unexamined Patent Application, First Publication No. 2005-105367    Patent Document 5: Japanese Unexamined Patent Application, First Publication No. 2007-302918    Patent Document 6: Japanese Unexamined Patent Application, First Publication No. 2006-52455    Patent Document 7: Japanese Unexamined Patent Application, First Publication No. S63-293121    Patent Document 8: Japanese Unexamined Patent Application, First Publication No. S57-137453    Patent Document 9: Japanese Unexamined Patent Application, First Publication No. 2001-152287    Non-Patent Document 1: Nissan Technical Review, No. 57 (2005-9), p. 4    Non-Patent Document 2: CAMP-ISIJ vol. 13 (2000), p. 411    Non-Patent Document 3: CAMP-ISIJ vol. 5 (1992), p. 1839    Non-Patent Document 4: CAMP-ISIJ vol. 13 (2000), p. 391